Steel part having long rolling contact fatigue life and method for producing the same

ABSTRACT

A steel part having a long rolling contact fatigue life and capable of further increasing the life of a bearing under severer using condition than usual conditions. The steel part includes steel having a composition containing 0.7% by mass to 1.1% by mass of C, 0.5% by mass to 2.0% by mass of Si, 0.4% by mass to 2.5% by mass of Mn, 1.6% by mass to 5.0% by mass of Cr, 0.1% by mass to less than 0.5% by mass of Mo, 0.010% by mass to 0.050% by mass of Al, less than 0.0015% by mass of Sb as an impurity, and the balance composed of Fe and inevitable impurities, the steel being hardened and tempered. In the steel structure of a portion from the surface to a depth of 5 mm, residual cementite has a grain diameter of 0.05 to 1.5 μm, prior austenite has a grain diameter of 30 μm or less, and the ratio by volume of the residual austenite is less than 25%.

RELATED APPLICATION

This application claims priority of Japanese Patent Application No.2006-052688, filed Feb. 28, 2006, herein incorporated by reference.

TECHNICAL FIELD

The technology in this disclosure relates to steel parts which have along rolling contact fatigue life and which are used as components ofrolling bearings such as roller bearings and ball bearings, and toroidalcontinuously variable transmissions. In particular, it relates to steelparts which have a long life until peculiar damage has occurred in asevere environment of bearing using, i.e., until microstructural change(damage) has occurred below a rolling contact plane due to cyclic load,and a method for producing the steel parts.

BACKGROUND

As materials for steel parts constituting rolling bearings used forautomobiles and industrial machines, high-carbon chromium bearing steeldefined in JIS-SUJ2 is most frequently used. In general, an importantproperty of bearing steel is a long rolling contact fatigue life, and apossible main factor which influences the rolling contact fatigue lifeis a non-metallic inclusion in steel. Therefore, as a commonly employedcountermeasure, the oxygen content in the high-carbon chromium steel isdecreased to control the amount, shape, and size of a non-metallicinclusion, thereby improving a bearing life (refer to, for example,Japanese Unexamined Patent Application Publication No. 1-306542 andJapanese Unexamined Patent Application Publication No. 3-126839).

However, in order to produce bearing steel containing a small amount ofnon-metallic inclusion, it is necessary to install expensive refiningequipment or significantly improve conventional equipment. Therefore,there is the problem of a high economic load.

Accordingly, research was conducted to resolve the problem. As a result,it was found that even when the amount of a non-metallic conclusion issimply decreased, in many cases, a large effect cannot be obtained onimprovement in the rolling life of a bearing, particularly the bearinglife under a severe condition such as a high load or a high temperature.This led to the finding that as a factor which determines the rollinglife, there is a factor other than the presence of a “non-metallicinclusion” which has been conventionally discussed. Specifically, amicrostructural change layer composed of a white etched constituentoccurs in a lower layer (surface layer) of a contact plane due to shearstress in contact between inner and outer rings and a rolling element ofa bearing as the environment of bearing using becomes severe. Inaddition, the microstructural change layer is gradually grown as thenumber of cycles increases, and finally spalling occurs by rollingcontact fatigue in the microstructural change portion to determine thebearing life. It was also found that the severe environment of bearingusing, i.e., a higher plane pressure (reduction in size) and an elevatedusing temperature, decrease the number of cycles until a microstructuralchange has occurred, resulting in a significant decrease in the bearinglife. Such a decrease in the bearing life in the severe environment ofusing cannot be sufficiently suppressed only by controlling the amountof a non-metallic inclusion as in related art. Therefore, it is thoughtto be necessary to retard the microstructural change.

As a countermeasure, bearing steel containing 0.5 to 1.5% by mass of C,over 2.5 to 8.0% by mass of Cr, 0.001 to 0.015% by mass of Sb, 0.002% bymass or less of 0, and the balance composed of Fe and inevitableimpurities has been proposed, and bearing steel containing theseelements and further containing over 0.5 to 2.5% by mass of Si, 0.05 to2.0% by mass of Mn, 0.05 to 0.5% by mass of Mo, and 0.005 to 0.07% bymass of Al has been developed (refer to Japanese Unexamined PatentApplication Publication No. 6-287691).

As a result, the microstructural change due to cyclic load in rollingcontact under high load was retarded, and so-called “B₅₀ high-loadrolling contact fatigue life (total number of cycles until a whiteportion of a microstructural change layer spalls at a cumulative failureprobability of 50% in a rolling contact fatigue test)” was improved.

However, the environment of bearing using has been recently made severerthan that at the time of filing of Japanese Unexamined PatentApplication Publication No. 6-287691, and thus the development of steelhaving a long rolling contact fatigue life has been desired ardently.

Therefore, a steel was developed having a long rolling contact fatiguelife as steel capable of further increasing a bearing life even undersevere using conditions, the steel having a composition containing 0.7to 1.1% by mass of C, 0.5 to 2.0% by mass of Si, 0.4 to 2.5% by mass ofMn, 1.6 to 4.0% by mass of Cr, 0.1 to less than 0.5% by mass of Mo,0.010 to 0.050% by mass of Al, and the balance composed of Fe andinevitable impurities, being subjected to hardening and tempering, andhaving a microstructure including residual cementite with a graindiameter of 0.05 to 1.5 μm and prior austenite with a grain diameter of30 μm or less (refer to Japanese Unexamined Patent ApplicationPublication No. 2004-315890).

With that steel, the average grain diameter of residual cementite isproperly controlled to retard the microstructural change and increasethe n number of cycles until the microstructure spalls. Furthermore, thegrain diameter of prior austenite in the microstructure after hardeningand tempering is refined to suppress the development of fatigue crackingand further improve the rolling contact fatigue life.

However, when the steel is applied to a component of an actual bearing,a sufficient rolling contact fatigue life may not be exhibited, therebycausing the need for further improvement.

SUMMARY

We provide steel parts which have a long rolling contact fatigue lifeand which are capable of further increasing the life of a bearing underseverer using conditions than usual conditions, and provide a usefulmethod for producing the steel parts.

We provide in particular:

-   -   (1) A steel part having a long rolling contact fatigue life,        comprising steel having a composition containing:        -   C: 0.7% by mass to 1.1% by mass;        -   Si: 0.5% by mass to 2.0% by mass;        -   Mn: 0.4% by mass to 2.5% by mass;        -   Cr: 1.6% by mass to 5.0% by mass;        -   Mo: 0.1% by mass to less than 0.5% by mass;        -   Al: 0.010% by mass to 0.050% by mass;    -    less than 0.0015% by mass of Sb as an impurity, and the balance        composed of Fe and inevitable impurities, the steel being        hardened and tempered, wherein a portion from the surface to a        depth of 5 mm has a steel structure in which residual cementite        has a grain diameter of 0.05 to 1.5 μm, prior austenite has a        grain diameter of 30 μm or less, and the ratio by volume of the        residual austenite is less than 25%.    -   (2) The steel part having a long rolling contact fatigue life        described above in (1), the steel further containing at least        one selected from the following:        -   Ni: 0.5% by mass to 2.0% by mass;        -   V: 0.05% by mass to 1.00% by mass; and        -   Nb: 0.005% by mass to 0.50% by mass.    -   (3) A method for producing a steel part having a long rolling        contact fatigue life, the method including hot-working steel,        spheroidizing annealing the steel by maintaining at 800° C. to        850° C. for 5 hours or more and cooling to 700° C. or less at a        rate of 0.01° C./s or less, and hardening and tempering the        steel, the steel having a composition containing:        -   C: 0.7% by mass to 1.1% by mass;        -   Si: 0.5% by mass to 2.0% by mass;        -   Mn: 0.4% by mass to 2.5% by mass;        -   Cr: 1.6% by mass to 5.0% by mass;        -   Mo: 0.1% by mass to less than 0.5% by mass;        -   Al: 0.010% by mass to 0.050% by mass;    -    less than 0.0015% by mass of Sb as an impurity, and the balance        composed of Fe and inevitable impurities.    -   (4) A method for producing a steel part having a long rolling        contact fatigue life, the method including hot-working steel,        cooling the steel to 200° C. at a cooling rate of 0.5° C./s or        less, spheroidizing annealing the steel by maintaining at        750° C. to 850° C. and cooling to 700° C. or less at a rate of        0.015° C./s or less, and hardening and tempering the steel, the        steel having a composition containing:        -   C: 0.7% by mass to 1.1% by mass;        -   Si: 0.5% by mass to 2.0% by mass;        -   Mn: 0.4% by mass to 2.5% by mass;        -   Cr: 1.6% by mass to 5.0% by mass;        -   Mo: 0.1% by mass to less than 0.5% by mass;        -   Al: 0.010% by mass to 0.050% by mass;    -    less than 0.0015% by mass of Sb as an impurity, and the balance        composed of Fe and inevitable impurities.    -   (5) The method for producing the steel part having a long        rolling contact fatigue life described above in (3) or (4), the        steel further containing at least one selected from the        following:        -   Ni: 0.5% by mass to 2.0% by mass;        -   V: 0.05% by mass to 1.00% by mass; and        -   Nb: 0.005% by mass to 0.50% by mass.

According to the present invention, a microstructural change in arolling environment under a high load is retarded, and thus a steel parthaving a high-load rolling contact fatigue life represented by so-calledB₅₀ can be provided. Therefore, a steel part required to have theresistance to rolling contact fatigue as a constituent part of, forexample, a roller bearing, can be reduced in size, and a steel partusable in an environment at a higher speed and higher load can beprovided.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a partial sectional view showing a toroidal continuouslyvariable transmission.

FIG. 2 is a schematic drawing showing a durability life test rig used ina rolling contact fatigue test.

In the drawings, reference numerals denote the following:

-   -   5 input disk    -   5 b orbital plane    -   6 b output disk    -   6 c orbital plane    -   7 roller    -   7 b peripheral surface    -   10 durability life test rig    -   11 disk    -   13 first roller    -   14 second roller    -   15 drive unit

DETAILED DESCRIPTION

We researched damage to a microstructure of steel used for a bearingunder a severe rolling environment. As a result, it was found that thedamage is mainly caused by stress concentration in a hard portion ofsteel and diffusion of carbon (symbol: C) in the periphery thereof. Inother words, a microstructural change in the steel can be retarded bysuppressing C diffusion in the steel under a using environment.

Therefore, a method for realizing the finding was further researched. Asa result, it was found that a method for suppressing C diffusion insteel, the austenite grains (represented) by γ hereinafter) present inthe metallic structure of the steel are refined in a heating process forhardening, and the grain diameter of retained cementite after hardeningand tempering is controlled to 0.05 to 1.5 μm.

In high-carbon bearing steel represented by JIS-SUJ2, coarse carbidewith a grain diameter of 5 μm or more, which is referred to as “eutecticcarbide”, may remain in steel after hardening and tempering due to theinfluence of coarse carbide crystallized when melted steel is cast andsolidified. Such coarse carbide is removed of course, and the spheroidalcarbide produced in spheroidizing annealing functions as a stressconcentrator with coarsening of the spheroidal carbide to promote amicrostructural change. As a measure against this, the steel disclosedin Japanese Unexamined Patent Application Publication No. 2004-315890was developed.

In other words, in the technique disclosed in Japanese Unexamined PatentApplication Publication No. 2004-315890, in order to retard amicrostructural change, the average grain diameter of residual cementiteis controlled in a proper range, specifically 0.05 to 1.5 μm, in whichstress concentration in a boundary between the residual cementite and amatrix can be suppressed while promoting dissolution of C into thematrix. In addition, as a method for controlling the average graindiameter of residual cementite in the proper range, a method ofspheroidizing annealing by maintaining at 750° C. to 850° C. and thencooling to 700° C. or less at a rate of 0.015° C./s or less is used.

Furthermore, in an attempt to apply the steel to an actual part, weworked the steel into the part shape by hot-working such as hot-castingor the like and then spheroidized and annealed the part under theconditions descried above to produce a desired microstructure. In thisattempt, the average grain diameter of residual cementite in a surfacelayer was not necessarily controlled in the range of 0.05 to 1.5 μm,thereby causing the problem of failing to obtain an expected rollingcontact fatigue characteristic. Furthermore, when finishing wasperformed by cutting after spheroidizing annealing (before hardening andtempering), the problem of very low machinability due to the hardsurface layer occurred.

As a result of intensive research on the causes of the problems, it wasfound that when large amounts of Cr and Mo are contained, themicrostructure of the surface layer after hot-working becomes astructure containing bainite or martensite, not containing cementite,and thus the growth of cementite does not sufficiently proceed even bysubsequent spheroidizing annealing. Since the growth of cementite isinsufficient after spheroidizing annealing, dissolution of C in anaustenite phase as a mother phase excessively proceeds during heatingfor hardening, and thus the amount of residual austenite after hardeningand tempering is excessively increased, thereby causing an adverseeffect on the rolling contact fatigue life. It was further found thatsince the microstructure of the surface layer becomes a structurecontaining bainite or martensite after hot-working, softening does notsufficiently proceed after spheroidizing annealing, thereby causingdifficulty in cutting.

As a result of further research, it was found that when hot-workingconditions or spheroidizing annealing conditions are optimized, themicrostructure of the surface layer of the steel part can be optimized,and the steel part having a long rolling contact fatigue life whilemaintaining machinability after spheroidizing annealing can be obtainedafter hardening and tempering.

Examples of the part include orbital parts such as an inner ring and anouter ring, and a rolling element, which constitute a rolling bearing,and a disk and a roller which constitute a toroidal continuouslyvariable transmission. FIG. 1 shows the structure of a toroidalcontinuously variable transmission as an example.

FIG. 1 is a schematic drawing showing a variator 1 of a full toroidalcontinuously variable transmission which is a type of toroidalcontinuously variable transmission. The variator 1 includes an inputshaft 3 rotated by an output shaft 2 of an engine, input disks 5 beingsupported near both ends of the input shaft 3.

In each of the input disks 5, a concavely curved orbital plane 5 b isformed in one of the sides, and a plurality of spline holes 5 a isformed in the inner periphery. The spline holes 5 a are engaged to asplined shaft 3 a provided on the input shaft 3 to integrally rotatablyattach each of the input disks 5 to the input shaft 3. In addition, theopposite movements of the input disks 5 are restricted by anchor rings51 fixed to the input shaft 3.

Further, output units 6 each including an output part 6 a and an outputdisk 6 b integrally rotatably supported by the output part 6 a arerelatively rotatably provided at the center of the input shaft 3 in theaxial direction thereof. In addition, a concavely curved orbital plane 6c is formed on one of the sides of each output disk 6 b which faces theorbital plane 5 b of each input disk 5. Further, sprocket gears 6 e areformed in the outer periphery of each output part 6 a so as to engagewith a chain 6 d so that power is transmitted to the outside through thechain 6 d.

Each of the output disks 6 b is attached to allow slight movement in theaxial direction of the output part 6 a, and a backup plate 6 h isdisposed at the back of each output disk 6 b with a gap 6 gtherebetween. The gap 6 g is sealed with a casing 6 f and a seal (notshown in the drawing). When hydraulic pressure is supplied to the gap 6g from a hydraulic power source, the output disk 6 b is urged toward theopposing input disk 5 to apply a predetermined terminal load.

In addition, a toroidal gap is formed between the orbital plane 5 b ofeach input disk 5 and the orbital plane 6 c of each output disk 6 b,which are opposed to each other. A lubricant (traction oil) is suppliedto the toroidal gap, and three disk rollers 7 are disposed at peripheralpositions at equal intervals so as to rotate in contact with the orbitalplanes 5 b and 6 c through an oil film. The portions of contact betweenthe orbital planes 5 b and 6 c are the peripheral surfaces 7 bb of therollers 7. Each of the rollers 7 is rotatably supported by a carriage 8so that the rotational shaft 7 a thereof can be tilted. In addition,driving forces is applied to the carriages 8 by hydraulic pressure inthe direction crossing the drawing of FIG. 1.

In the variator 1, when the pair of input disks 5 is rotated, torque istransmitted from the input disks 5 to the output disks 6 b by shearingforce of the oil film through the three rollers 7 on each of the rightand left sides. The rollers 7 supported by the carriages 8 incline therotational shafts 7 a to remove unbalance between the reaction forcegenerated in the carriages 8 due to torque transmission and the torquenecessary for driving the output disks 6 b. Consequently, the positionsof the rollers 7 are changed as shown by two-dot chain lines in FIG. 1to continuously change the change gear ratio between the disks 5 and 6b.

In such a toroidal continuously variable transmission, the portions ofrolling contact between the input/output disks and the rollers aresubjected to a high temperature (100° C. or more) and high surfacepressure (maximum contact surface pressure 4 to 4.5 GPa or more). Inaddition, for example, when three rollers are disposed, large verticalstress is repeatedly applied to the orbital planes of the input/outputdisks from the three points, and high shearing stress is repeatedlyapplied due to the traction during power transmission. Therefore, theorbital planes of the input/output disks are under specific severecontact conditions as compared with a rolling surface of a usual rollingbearing in which only vertical stress is mainly applied to the surface.Therefore, the orbital plane of each disk is required to have highfatigue strength in order to prevent the occurrence of spalling startingat the surface due to metal contact and spalling due to a structuralchange even when used under such specific severe contact conditions.

Next, the reasons for providing the ranges of the components of thesteel part will be described below.

C: about 0.7% by Mass to about 1.1% by Mass

In a metal structure of steel, C is dissolved in a matrix. However, inorder to strengthen martensite grains to cause an effective function tosecure hardness of steel after hardening and tempering and improve therolling contact fatigue life, a C content of about 0.7% by mass or moreis required. However, at an excessively high C content, the formation ofcoarse carbide such as eutectic carbide is accelerated, and amicrostructural change due to C diffusion in steel is promoted, therebydegrading the rolling contact fatigue life. Therefore, the upper limitis about 1.1% by mass.

Si: about 0.5% by Mass to about 2.0% by Mass

Si functions as a deoxidizer in refining of steel and is solid-dissolvedin a matrix to suppress a decrease in strength of steel in temperingafter hardening. Further, Si is an effective element for retarding amicrostructural change in an environment under rolling load. In order tosufficiently exhibit these effects, a Si content of about 0.5% by massor more is required. On the other hand, a content over about 2.0% bymass causes significant deterioration in the castability andmachinability of steel. Therefore, the upper limit is about 2.0% bymass.

Mn: about 0.4% by Mass to about 2.5% by Mass

Mn functions as a deoxidizer in refining of steel and is an effectiveelement for decreasing the oxygen content of steel. In addition, Mneffectively functions to improve the hardenability of steel to improvetoughness and strength of martensite constituting a matrix, and improvethe rolling contact fatigue life. Furthermore, Mn has the effect ofstabilizing a cementite phase and retard a microstructural change. Inorder to sufficiently exhibit these effects, a Mn content of about 0.4%by mass or more is required. On the other hand, a content over about2.5% by mass causes significant deterioration in the castability andmachinability of steel. Therefore, the upper limit is about 2.5% bymass.

Cr: about 1.6% by Mass to about 5.0% by Mass

Cr has the function to stabilize a cementite phase in steel to suppressC diffusion and the function to suppress coarsening of cementite grainsto prevent stress concentration. Also, Cr is an element effectivelyfunctioning to improve the rolling contact fatigue life. In order toachieve the sufficient effect, a Cr content of about 1.6% by mass ormore is required. On the other hand, at a content of over about 5% bymass, the amount of C dissolved in martensite is decreased to decreasehardness after hardening and tempering, thereby degrading the rollingcontact fatigue life. Therefore, the upper limit is about 5% by mass.

Mo: about 0.1% by Mass to Less than about 0.5% by Mass

Mo is solid-dissolved in a matrix and has the function to suppress adecrease in strength of steel during tempering after hardening. Further,Mo improves hardness of steel after hardening and tempering and improvesthe rolling contact fatigue life. In addition, Mo has the function tostabilize carbide to retard a microstructural change. However, at acontent of less than about 0.1% by mass, the effect cannot besufficiently obtained, while at a content of about 0.5% by mass or more,the effect is saturated, thereby increasing the cost. Therefore, thecontent is about 0.1% by mass to less than about 0.5% by mass.

Al: about 0.010% by Mass to about 0.050% by Mass

Al is necessary as a deoxidizer in refining of steel and is positivelyadded as an element functioning to refine prior austenite grains bybonding to N in steel to effectively improve the rolling contact fatiguelife. In order to obtain the sufficient effect, a content of about0.010% by mass or more is required. On the other hand, at a high contentover about 0.050% by mass, the rolling contact fatigue life is degradedby AlN precipitated in a large amount in steel. Therefore, the contentis about 0.010% by mass to about 0.050% by mass.

Sb: Less than about 0.0015% by Mass

Sb is an element possibly mixed from an iron source such as scraps, butsegregates at austenite grain boundaries in hot-working to degradehot-workability, toughness, and the rolling contact fatigue life ofsteel due to mixing of Sb. Therefore, it is necessary to control theamount of Sb mixing to a low value by appropriately selecting an ironsource. The above problem generally becomes significant when about0.0010% by mass or more Sb is mixed. However, when the grain diameter ofprior austenite is achieved, the allowable upper limit of the amount ofSb mixing can be increased by increasing the grain boundary area.However, the amount of Sb mixing in steel must be controlled to lessthan about 0.0015% by mass.

In addition to the above-described basic components, at least oneselected from about 0.5% by mass to about 2.0% by mass of Ni, about0.05% by mass to about 1.00% by mass of V, and about 0.005% by mass toabout 0.50% by mass of Nb can be further contained.

Ni: about 0.5% by Mass to about 2.0% by Mass

Since Ni is solid-dissolved in a matrix to suppress a decrease instrength of steel after tempering, Ni is added according to demand. Inorder to obtain the sufficient effect, a content of about 0.5% by massis required. On the other hand, a content of over about 2.0% by masscauses the formation of a large amount of residual austenite, therebydecreasing the strength after hardening and tempering. Therefore, theupper limit of the content is about 2.0% by mass.

V: about 0.05% by Mass to about 1.00% by Mass

V has the function to form stable carbide and improve hardness of steeland the function to suppress a microstructural change to improve therolling contact fatigue life. Therefore, V is added according to demand.In this case, at a content of less than about 0.05% by mass, thesufficient effect is not obtained, while at an excessively high content,the amount of dissolved C is decreased to decrease hardness of steelafter hardening and tempering. Therefore, the upper limit of the contentis about 1.00% by mass.

Nb: about 0.005% by Mass to about 0.50% by Mass

Like V, Nb has the function to form stable carbide and improve hardnessof steel and the function to suppress a microstructural change toimprove the rolling contact fatigue life. Therefore, Nb is addedaccording to demand. In this case, at a content of less than about0.005% by mass, the sufficient effect cannot be obtained, while at acontent over about 0.05% by mass, the effect is saturated. Therefore,the content is about 0.005% by mass to about 0.50% by mass.

Next, the microstructure of the steel part will be described below.

It was found that in a steel part required to have a rolling contactfatigue life, the microstructure of a surface layer from the surface toa depth of about 5 mm is particularly important. Therefore, the steelpart after hardening and tempering is required to have a portion fromthe surface to a depth of about 5 mm which satisfies the microstructuredescribed below.

First, in the steel part after hardening and tempering, the cementitegrain diameter in the steel structure of a portion from the surface to adepth of about 5 mm is controlled to about 0.05 μm to about 1.5 μm forthe following reasons:

-   -   When the steel having the above-described C content is hardened        and tempered, cementite present before hardening remains in the        microstructure of the steel. Therefore, we repeatedly studied        with attention to the fact that a distribution form of the        residual cementite strongly influences the properties of a        microstructural change. As a result, it was found that when the        average grain diameter of residual cementite is smaller than        about 0.05 μm, the ratio of the surface area of cementite to the        volume is increased to promote dissolution of C into the matrix.        On the other hand, when coarse residual cementite having an        average grain diameter over about 1.5 μm is present, stress        concentration in the boundaries between the residual cementite        and the matrix is accelerated, and thus the number of cycles of        stress application until the occurrence of a microstructural        change and spalling of the microstructure is decreased. From        this viewpoint, it was found that the grain diameter of residual        cementite after hardening and tempering is preferably specified        to about 0.05 μm to about 1.5 μm.    -   In the steel structure of a portion from the surface to a depth        of about 5 mm, the grain diameter of prior austenite is        specified to about 30 μm or less. The reason for this is that in        the microstructure of the steel after hardening and tempering,        when the grain diameter of prior austenite is about 30 μm or        less, propagation of fatigue cracks within a crystal grain can        be stopped at the grain boundary, and further progress can be        retarded.    -   Furthermore, in the steel structure of a portion from the        surface to a depth of about 5 mm, the ratio by volume of the        residual austenite is specified to less than about 25%. In other        words, a residual austenite phase at a residual austenite ratio        of about 25% or more is transformed to martensite accompanying        volume expansion in the using environment, thereby changing the        dimensions of the steel part. The dimensional change causes a        stress concentration portion to adversely affect the rolling        contact fatigue.

The steel part having a long rolling contact fatigue life is producedthrough the following steps:

-   -   First, molten steel having the above-described chemical        composition is refined in a steel making process and then        continuously cast to form a cast slab. The steel cast slab is        formed into a steel material (for example, a steel bar) by a        hot-rolling process. The steel material is then formed in a        steel part such as a bearing race by hot-working such as hot        casting or the like. After spheroidizing annealing, if required,        the steel part is cut and hardened and tempered to produce a        steel part.

In the method for producing the steel part, it is necessary to use thefollowing conditions (I) or (II):

-   -   (I) After the hot working, the steel is spheroidized and        annealed by maintaining at about 800° C. to about 850° C. for        about 5 hours or more and then cooling at a rate of about 0.01°        C./s or less, and then hardened and tempered.    -   (II) After hot working at abut 900° C. or more, the steel is        cooled to about 200° C. at a cooling rate of about 0.5° C./s or        less, spheroidized and annealed by maintaining at about 750° C.        to about 850° C. and then cooling to about 700° C. at a cooling        rate of about 0.015° C./s or less, and then hardened and        tempered.

In the method (I), the spheroidizing annealing conditions are controlledto control the average grain diameter of cementite in the surface layerof the steel part. When the usual hot-working conditions and subsequentcooling conditions are not particularly limited, in the steelcomposition, the structure of the surface layer may become a bainite ormartensite structure after hot-working and cooling. Therefore, it isnecessary to produce cementite for achieving the above-described finalstructure by subsequent spheroidizing annealing of the steel having abainite or martensite structure. As a condition for this purpose, aspheroidizing annealing condition is required, in which the steel ismaintained at about 800° C. to about 850° C. for about 5 hours or moreand then cooled to about 700° C. or less at a cooling rate of about0.01° C./s or less.

Namely, when the retention temperature is lower than about 800° C. orthe retention time is less than about 5 hours, the sufficient growth ofcementite in the surface layer cannot be expected, and the cementitegrain diameter in the final surface layer structure cannot be controlledin the above-described range. Also, solid-dissolution of C in theaustenite phase, which is a matrix in hardening heating, excessivelyproceeds, and thus the amount of the residual austenite in the finalsurface layer structure is excessively increased. On the other hand,when the retention temperature exceeds about 850° C., cementite afterspheroidizing is coarsened, and, consequently, the residual cementiteafter hardening and tempering is also coarsened. In cooling after theretention, when the cooling rate to about 700° C. exceeds about 0.01°C./s, cementite precipitation during cooling proceeds in the form ofreproduction of pearlite, not the growth of spheroidized cementite.Therefore, softening after spheroidizing annealing does not sufficientlyproceed to degrade workability. Further, solid dissolution in hardeningheating is excessively accelerated to form a large amount of residualaustenite after hardening.

After the above-described spheroidizing annealing, hardening andtempering is performed. In order to obtain a desired residual cementitedistribution and prior austenite grain diameter after hardening andtempering, the heating temperature of hardening is preferably about 800°C. to about 950° C. This is because the most desirable microstructure isproduced in this temperature range. Although the fraction of thecementite structure after hardening and tempering changes mainlydepending on the C content, the volume ratio is about 3 to about 25% inthe composition range of the present invention. In addition, a cuttingwork may be performed before the hardening and tempering. Theabove-mentioned spheroidizing annealing conditions have the effect ofimproving machinability because the surface layer is sufficientlysoftened.

Next, the method (II) will be described.

In the method (II), the hot-working conditions and subsequent coolingconditions are regulated to control the surface layer structure afterthe hot-working and cooling. When the hot-working temperature is lowerthan about 900° C., the mold life is decreased due to the highdeformation resistance of steel, and cracks occur in casting due to thelow deformation ability. Therefore, the hot-working is performed atabout 900° C. or more. Further, in subsequent cooling, the condition iscontrolled so that the above-described final structure can be obtainedin the surface layer.

When the cooling rate exceeds about 0.5° C./s, the sufficient growth ofcementite in the surface layer cannot be expected by subsequentspheroidizing annealing, and the cementite grain diameter in the finalsurface layer structure cannot be controlled in the above-describedrange. Also, solid-dissolution of C in the austenite phase, which is amother phase in hardening heating, excessively proceeds, and thus theamount of the residual austenite in the final surface layer structure isexcessively increased.

Then, spheroidizing annealing is performed by retention at about 750° C.to about 850° C. and then cooling to about 700° C. at a cooling rate ofabout 0.015° C./s or less. When the retention temperature exceeds about850° C., care must be taken because cementite after spheroidizing iscoarsened, and, consequently, the residual cementite after hardening andtempering is also coarsened. At the same time, layered cementite isnewly formed in cooing after the retention, thereby causing difficult inobtaining desired spheroidal cementite.

On the other hand, when the retention temperature is lower than about750° C., decomposition of cementite present as pearlite beforespheroidizing annealing does not sufficiently proceed, and thus adesired residual cementite distribution cannot be obtained.

In cooling after the retention, the cooling rate to about 700° C. orless must be about 0.015° C./s or less. When the cooling rate exceedsabout 0.015° C./s, cementite precipitation during cooling proceeds inthe form of reproduction of pearlite, not the growth of spheroidizedcementite. Therefore, softening after spheroidizing annealing does notsufficiently proceed to degrade workability. Further, solid dissolutionin hardening heating is excessively accelerated to form a large amountof residual austenite after hardening.

After the above-described spheroidizing annealing, hardening andtempering is performed. In order to obtain a desired residual cementitedistribution and prior austenite grain diameter after hardening andtempering, the heating temperature of hardening is preferably about 800°C. to about 950° C. This is because the most desirable microstructure isproduced in this temperature range. Although the fraction of thecementite structure after hardening and tempering changes mainlydepending on the C content, the volume ratio is about 3 to about 25% inthe composition range of the present invention. In addition, a cuttingwork may be performed before the hardening and tempering. Theabove-mentioned spheroidizing annealing conditions have the effect ofimproving machinability because the surface layer is sufficientlysoftened.

EXAMPLES

Molten steel having each of the chemical compositions shown in Table 1was refined by a converter and then continuously cast to form a caseslab. The resulting cast slab was diffusion-annealed at 1200° C. for 30hours and then rolled to a steel bar of 64 mm or 90 mm in diameter.

The steel bar of 64 mm in diameter and the steel bar of 90 mm indiameter were hot-cast into a disk-like roller shape and a disk shape,respectively, at least the temperature shown in Table 2 and then cooledat various cooling rates. These disk and roller were normalized and thenspheroidized and annealed. The spheroidizing annealing was performed bycooling to 650° C. from various retention temperatures at variouscooling rates shown in Table 2 and then standing to cool. Then, in orderto remove a decarbonized layer, a cutting work was performed to form atest piece with a final shape.

Further, after hardening, the tempering temperature was changed from theheating temperature shown in Table 2 according to the steel used, andthe hardness HRc after tempering was controlled to 60 to 62, followed bypolishing and lapping finishing.

Next, the resulting test piece was cut in the height direction of acolumn, and a section was corroded with a picric acid alcohol solutionand then corroded with a nitric acid alcohol solution. Then, themicrostructure was observed to measure the average grain diameter ofresidual cementite and the average grain diameter of prior austenite byimage analysis.

A rolling contact fatigue test was conducted using a durability lifetest rig 10 of a traction transmission part shown in FIG. 2. Thedurability life test rig shown in FIG. 2 includes a disk 11, a disksupport 12 for supporting the disk 11, a first roller 13 to be inrolling contact with one of the sides of the disk 11, a second roller 14to be in rolling contact with the other side of the disk 11, a driveunit 15 for rotating the first roller 13, a differential rate mechanism16 giving a peripheral speed difference to the second roller 14 relativeto the first roller 13, and a pressure unit 17 for pressing the firstroller 13 and the second roller 14 on the disk 11, all of which aredisposed on a pedestal A.

In the durability life test rig having the above-described constitution,in the state in which the disk 11 is supported by the disk support 12,and the peripheral side of the disk 11 is held between the first roller13 and the second roller 14, the disk can be caused to follow therollers 13 and 14 by the pressure unit 17 to cause rolling contact eachof the rollers 13 and 14 and the disk 11. In this case, a peripheralspeed difference can be given to the second roller 14 by thedifferential speed mechanism 16 relative to the first roller 13.Therefore, slipping can be caused between the disk 11 and the secondroller 14.

The durability life test can be conducted by reproducing a use state inan actual traction transmission part. In the durability life test rig,the above-described test piece was used for the disk 11, the firstroller 13, and the second roller 14. The time taken until failure hadoccurred due to rolling slide with the roller 14 was measured as adurability life, and the durability of the disk 11 was evaluated on thebasis of the durability life.

The test conditions of the durability life test rig 10 were determinedas follows:

-   -   (1) Roller rotational speed: 3000 rpm    -   (2) Slip ratio between roller and disk: 14%    -   (3) Maximum contact surface pressure: 4.2 GPa    -   (4) Lubricant: traction oil for toroidal continuously variable        transmission    -   (5) Oil film parameter (Λ): 1.8        The results of the durability life test are shown in Table 2. In        Table 1, steel No. 1 is conventional steel corresponding to        JIS-SUJ2, and steel Nos. 2, 3, 4, 5, 6, 6, and 8 are comparative        steel containing C, Si, Mn, Cr, Mo, Al, and Sb, respectively, at        contents out of the range of the present invention.

In steel part Nos. 2 to 8 (comparative steel) containing essentialelements at contents out of the specified range, B₅₀ is inferior tosteel part No. 1 but is substantially the same value as steel No. 1. Inparticular, in steel part Nos. 2 and 7 using steels having a low Ccontent and a low Al content, respectively, the microstructure cannot becontrolled in the specified range, and the value of B₅₀ is excessivelydecreased. In steel part Nos. 9, 11, 20, and 21 each using steel inwhich the chemical composition is within the range, but themicrostructure differs from the specified structure, B₅₀ is superior tosteel part No. 1 (conventional steel), but the improvement in B₅₀ isonly small.

On the other hand, in steel part Nos. 10, 12 to 19, and 21 in which boththe chemical composition and the microstructure are within the specifiedranges, B₅₀ is 10 times or more superior to steel part No. 1(conventional steel). These test results indicate that a damagepreventing effect is particularly large in a full-toroidal continuouslyvariable transmission having a large spin component in a contact portionbetween the roller and the disk and severe contact conditions.

TABLE 1 Steel Chemical composition of steel (% by mass) No. C Si Mn CrMo Al Ni V Nb Sb Remarks 1 1.00 0.25 0.50 1.50 0.01 0.033 — — — 0.0014Conventional steel 2 0.65 0.98 0.59 4.22 0.44 0.033 — — — 0.0008Comparative steel 3 0.90 0.13 0.68 2.51 0.22 0.030 — — — 0.0009 ″ 4 1.071.47 0.30 2.02 0.54 0.028 — — — 0.0008 ″ 5 0.86 1.22 0.53 1.20 0.510.033 — — — 0.0009 ″ 6 0.84 1.54 0.57 4.34 0.03 0.028 — — — 0.0010 ″ 70.88 1.04 0.69 2.65 0.31 0.005 — — — 0.0011 ″ 8 1.02 1.65 0.63 2.34 0.380.033 — — — 0.0038 ″ 9 1.00 1.00 0.45 3.50 0.45 0.035 — — — 0.0010 Steelof this invention 10 1.00 1.50 0.70 5.00 0.45 0.033 — — — 0.0011 ″ 110.95 1.10 1.50 2.00 0.30 0.027 — — — 0.0009 ″ 12 0.91 1.23 0.70 4.250.40 0.026 1.03 — — 0.0009 ″ 13 0.90 0.91 0.70 2.55 0.50 0.033 — 0.31 —0.0011 ″ 14 0.85 1.28 0.78 4.49 0.56 0.029 — — 0.045 0.0009 ″ 15 1.030.73 0.58 3.87 0.28 0.028 0.70 — 0.031 0.0010 ″ 16 1.07 0.74 0.62 4.340.25 0.032 0.62 0.15 — 0.0008 ″

TABLE 2 Hot-working Spheroidizing Surface layer structure after(casting) condition annealing condition hardening and tempering CoolingCooing Average Hot- rate after rate Hardening Average prior γ Amount ofSteel working hot- Retention Retention (retention Retention residual θgrain residual part Steel temp. working temp. time temp. ~700° C.) temp.diameter diameter γ (% by B₅₀ No. No. (° C.) (° C./s) (° C.) (h) (°C./s) (° C.) (μm) (μm) mass) life ratio Remarks 1 1 1000 0.4 810 8 0.007900 0.47 11.1 13.1 1.0 *1) 2 2 1100 0.4 810 8 0.007 900 0.02 14.8 10.90.5 *2) 3 3 1050 0.4 810 8 0.007 900 0.30 11.5 12.3 0.9 ″ 4 4 1000 0.4810 8 0.007 900 0.40 13.4 11.0 1.1 ″ 5 5 1050 0.4 810 8 0.007 900 0.3411.0 12.6 1.0 ″ 6 6 1050 0.4 810 8 0.007 900 0.30 13.5 12.5 1.2 ″ 7 71050 0.4 810 8 0.007 900 0.34 36.2 10.3 0.4 ″ 8 8 1000 0.4 820 7 0.007900 0.30 14.4 14.0 1.3 ″ 9 9 1000 1.0 780 7 0.008 900 0.03 19.3 30.0 2.6″ 10 9 1000 1.0 830 7 0.006 900 0.38 13.2 12.0 12.6 *3) 11 9 1000 1.0830 7 0.006 1000 0.06 45.0 24.0 2.3 *2) 12 9 1000 0.2 790 7 0.006 9000.20 11.6 15.0 10.8 *1) 13 10 1000 0.4 810 8 0.006 950 0.26 16.4 12.914.9 ″ 14 11 1000 0.4 820 7 0.007 900 0.43 13.0 13.6 3.3 ″ 15 12 10500.3 830 8 0.005 950 0.25 18.3 11.6 16.4 ″ 16 13 1050 0.2 780 8 0.005 9000.31 12.5 11.6 9.2 ″ 17 14 1050 2.0 845 7 0.003 950 0.25 17.6 11.4 16.9″ 18 15 1000 0.4 820 7 0.007 900 0.40 14.9 10.5 12.8 ″ 19 16 1000 2 7903 0.005 900 0.03 21.0 35.0 3.3 *2) 20 16 1000 1.5 830 8 0.052 900 0.0318.0 23.5 3.5 *2) 21 16 1000 1.5 830 8 0.005 900 0.35 13.5 11.5 14.4 *3)*1) Conventional example *2) Comparative example *3) Example of thisinvention

1. A steel part having a long rolling contact fatigue life, comprisingsteel having a composition containing: C: about 0.7% by mass to about1.1% by mass; Si: about 0.5% by mass to about 2.0% by mass; Mn: about0.4% by mass to about 2.5% by mass; Cr: about 1.6% by mass to about 5.0%by mass; Mo: about 0.1% by mass to less than about 0.5% by mass; Al:about 0.010% by mass to about 0.050% by mass; less than about 0.0015% bymass of Sb as an impurity, and the balance composed of Fe and inevitableimpurities, the steel being hardened and tempered, wherein a portionfrom the surface to a depth of about 5 mm has a steel structure in whichresidual cementite has a grain diameter of about 0.05 to about 1.5 μm,prior austenite has a grain diameter of about 30 μm or less, and theratio by volume of the residual austenite is less than about 25%.
 2. Thesteel part according to claim 1, wherein the steel further comprises atleast one selected from the following: Ni: about 0.5% by mass to about2.0% by mass; V: about 0.05% by mass to about 1.00% by mass; and Nb:about 0.005% by mass to about 0.50% by mass.
 3. A method for producing asteel part having a long rolling contact fatigue life, comprising:hot-working steel, spheroidizing annealing the steel by maintaining thesteel at about 800° C. to about 850° C. for about 5 hours or more,cooling the steel to about 700° C. or less at a rate of about 0.01° C./sor less, and hardening and tempering the steel, wherein the steel has acomposition comprising: C: about 0.7% by mass to about 1.1% by mass; Si:about 0.5% by mass to about 2.0% by mass; Mn: about 0.4% by mass toabout 2.5% by mass; Cr: about 1.6% by mass to about 5.0% by mass; Mo:about 0.1% by mass to less than about 0.5% by mass; Al: about 0.010% bymass to 0.050% by mass;  less than about 0.0015% by mass of Sb as animpurity, and the balance composed of Fe and inevitable impurities.
 4. Amethod for producing a steel part having a long rolling contact fatiguelife, comprising: hot-working steel, cooling the steel to about 200° C.at a cooling rate of about 0.5° C./s or less, spheroidizing annealingthe steel by maintaining at about 750° C. to about 850° C., cooling thesteel to about 700° C. or less at a rate of about 0.015° C./s or less,and hardening and tempering the steel, wherein the steel has acomposition comprising: C: about 0.7% by mass to about 1.1% by mass; Si:about 0.5% by mass to about 2.0% by mass; Mn: about 0.4% by mass toabout 2.5% by mass; Cr: about 1.6% by mass to about 5.0% by mass; Mo:about 0.1% by mass to less than about 0.5% by mass; Al: about 0.010% bymass to about 0.050% by mass;  less than about 0.0015% by mass of Sb asan impurity, and the balance composed of Fe and inevitable impurities.5. The method according to claim 3, wherein the steel further comprisesat least one selected from the following: Ni: about 0.5% by mass toabout 2.0% by mass; V: about 0.05% by mass to about 1.00% by mass; andNb: about 0.005% by mass to about 0.50% by mass.
 6. The method accordingto claim 4, wherein the steel further comprises at least one selectedfrom the following: Ni: about 0.5% by mass to about 2.0% by mass; V:about 0.05% by mass to about 1.00% by mass; and Nb: about 0.005% by massto about 0.50% by mass.